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Effect of sintering temperature on grain boundary character distribution in pure nickel
Scripta Materialia 56 (2007) 13–16 www.actamat-journals.com
Effect of sintering temperature on grain boundary character distribution in pure nickel P.P. Bhattacharjee, S.K. Sinha and A. Upadhyaya* Department of Materials and Metallurgical Engineering, Indian Institute of Technology, Kanpur 208 016, UP, India Received 25 May 2006; revised 29 August 2006; accepted 1 September 2006 Available online 10 October 2006
The evolution of grain boundaries in pure Ni has been studied after sintering at 900, 1100 and 1300 C. It is found that the fraction of low-angle grain boundaries remains more or less constant over this sintering temperature range. However, the fraction of random boundaries decreases and that of R3 boundaries increases with increasing sintering temperature. The observations on grain boundary character distribution are found to correlate well with the accompanying microstructural changes. 2006 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Nickel; Sintering; Grain growth; EBSD; Grain boundary
The phenomenon of grain growth following recrystallization has been well documented in the literature [1]. Grain growth has often been found to be associated with the development of a characteristic recrystallization texture and accompanying changes in grain boundary character distribution in a number of systems. The effect of grain growth and other annealing processes on grain boundary character distribution has been an area of significant interest due to the fact that grain boundaries can substantially influence properties such as creep resistance [2] and susceptibility to intergranular corrosion [3], and, more recently, critical current density in superconducting thin films has been found to be dependent on grain boundary character [4]. Enhancing the properties of materials by tailoring grain boundaries for a number of critical applications through ‘‘grain boundary engineering’’ has become an established and active area of materials research. However, grain growth and other annealing phenomena have generally been studied in great detail in materials that have been cold worked prior to the annealing treatment. Interestingly, in sintering, which is also a high-temperature treatment and is employed for densification of powder metallurgically prepared porous compacts, grain growth has been found to play a significant role. However, the evolution of grain boundaries * Corresponding author. Tel.: +91 512 2597672; fax: +91 512 2597505; e-mail: [email protected]
has been studied to a much lesser extent in these materials. This is surprising since powder metallurgically prepared materials are often of near net shape and hence there is a greater need to study the evolution of grain boundaries during sintering in such materials. With this in mind, the current work investigates the evolution of grain boundaries and the accompanying microstructural changes in pure Ni during sintering at different temperatures. Figure 1 shows a scanning electron microscope (SEM) micrograph while Table 1 summarizes the important characteristics of the Ni powder used in this study. It should be noted that the average particle diameter in Table 1 is evidently more than what appears in the micrograph of Figure 1. This may be because a wet measurement system (as in Table 1) usually gives larger particle diameters, due to particle agglomeration, compared to measurements on the population base (as in the SEM micrograph of Fig. 1). Pure Ni compacts were made using a rectangular die of dimensions 8 · 25 mm in a 50 ton hydraulic press (Model CTM 50, FIE, India) applying uniaxial pressure of 500 MPa. The green samples were sintered for 1 h at 900 C (±5C), 1100 C (±5 C) and 1300 C (±5 C), respectively. Sintering was carried out in a SiC-heated horizontal tubular furnace (rating 1.5 kV A) using commercially pure hydrogen as the sintering atmosphere (dew point 35 C). The heating rate (5 C min1) was kept constant for all the sintering operations. Green and sintered densities were determined from the dimensional
1359-6462/$ - see front matter 2006 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.scriptamat.2006.09.003
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P. P. Bhattacharjee et al. / Scripta Materialia 56 (2007) 13–16
Figure 1. SEM micrograph of the Ni powder.
Table 1. Characteristics of the experimental Ni powder Powder Vendor Purity Shape Particle size (lm)
D10 D50 D90
Ni Sigma–Aldrich 99.99% Spiky 3.9 9.7 26
measurements of the green and sintered compacts, respectively, and the densification parameter was calculated from these values using the relation: SD GD ; TD GD where SD is sintered density, GD is green density and TD is theoretical density. Specimens for optical microscopy studies were prepared using usual metallography techniques. Samples for orientation imaging microscopy (OIM) studies were prepared using electropolishing at room temperature with platinum as cathode and the samples as anode. The OIM studies were conducted by a fully computer-controlled electron backscattered diffraction (EBSD) system attached to a SEM (QUANTA 200, FEI) using TSL OIM ANALYSIS Version 4.0 software. The step sizes were chosen such that at least 3–4 Kikuchi patterns were collected from each grain. For statistical accuracy, approximately 2000 grains were covered for each sample. Table 2 gives the sintered densities and densification parameters of pure Ni sintered at three different temperatures. The sintered density is found to increase in a consistent manner with increasing sintering temperatures. Figure 2(a)–(c) shows the optical micrographs of pure Ni sintered at 900, 1100 and 1300 C, respectively. Annealing twins are clearly visible in the microstructures Densification parameter ¼
Table 2. Effect of sintering temperature on the densification behavior of pure Ni Sintering temperature, C
Sintered density, % theoretical
Densification parameter
900 1100 1300
82.0 90.0 94.0
0.47 0.68 0.73
Figure 2. Optical micrographs of Ni compacts sintered at (a) 900 C, (b) 1100 C and (c) 1300 C.
of the materials sintered at 1100 and 1300 C as compared to the material sintered at 900 C. The micrographs of materials sintered at 1100 and 1300 C clearly indicate signs of grain growth, a finding that is also substantiated by the grain size distributions (Fig. 3) for these three sintering temperatures. The average grain size was around 6 lm after sintering at 900 C and around 11 and 35 lm after sintering at 1100 and 1300 C, respectively. The grain boundary character distributions (GBCDs) of pure Ni at these three sintering conditions reveal that the fraction of low-angle grain boundaries (LAGBs) is quite low after sintering at 900 C and does not change to any significant extent after sintering at 1100 and 1300 C. However, the coincidence site lattice (CSL) and random grain boundary fractions differ markedly in these three conditions. The fraction of random grain boundaries is quite high after sintering at 900 C and decreases slightly after sintering at 1100 C and quite significantly after sintering at 1300 C.
P. P. Bhattacharjee et al. / Scripta Materialia 56 (2007) 13–16
Figure 3. Grain size distribution in pure Ni after sintering at (a) 900 C, (b) 1100 C and (c) 1300 C.
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boundaries in face-centered cubic (fcc) metals and alloys belong to the R3 (60/h1 1 1i relationship) category. The sintered microstructures and misorientation distribution plots also indicate that the R3 boundaries present can be identified exclusively as twin boundaries. The increase in R3 boundary fraction shows that the twin density increases with increasing sintering temperatures. In order to quantify the above observations, unique grain color maps were constructed from the EBSD datasets for these three sintered materials and volume fractions of the parent grains (P) or grains containing twins, the twinned regions within the parent grains (T) and the untwinned (U) grains (schematically shown in Fig. 5(a)) were determined (Fig. 5(b)) for the three sintering temperatures. Quantitative examination also shows clearly that the twin density increases with increasing sintering temperature. It has previously been suggested from experimental evidence [5–7] on different materials, as well as by computer simulations [8,9], that in a microstructure with a distribution of grain boundary energies the boundary tension would lead to a decrease in the total number of high-energy boundaries on annealing. In pure Ni, Furley and Randle [5] have reported an increase in R3 boundaries and a decrease in R5 boundaries during grain growth. However, in most cases the materials have been prepared through thermomechanical processing and as a result it may be difficult to pinpoint whether primary recrystallization or grain growth is the dominating effect in changing the GBCD. However, it can be seen that there is a general tendency for the proportion of low-energy grain boundaries to increase and high-energy grain boundaries to decrease during grain growth. Mobility of boundaries during grain growth has been thought to play a significant role, and a large increase is predicted for low-energy/low-mobility boundaries, such as LAGBs (R1) and other low R boundaries [9]. In the present study, pure Ni samples consolidated by the powder metallurgy technique have been investigated. Therefore, the complex effects of thermomechanical processing may be neglected and the change in GBCD can be solely attributed to changes taking place at the grain growth stage of sintering. The GBCDs clearly show that there is no significant change in LAGB proportion during grain growth. However, the fraction of random boundaries decreases and the proportion of R3 boundary increases with increasing
Figure 4. The CSL boundary distributions in pure Ni after sintering at (a) 900 C, (b) 1100 C and (c) 1300 C.
The CSL boundary fraction shows almost a reverse trend (Fig. 4(a)–(c)). The fraction of CSL boundary is rather low for the 900 C sintered sample and increases slightly after sintering at 1100 C and quite significantly after sintering at 1300 C. The CSL boundary distributions in these three conditions clearly indicate that most of the CSL boundaries belong to the R3 category. Boundaries beyond R3 are present in very minor proportions in all the sintered materials. Annealing twin
Figure 5. (a) Schematic representation of P, T and U (see text) and (b) histogram showing their respective volume percentages in the three sintered conditions.
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P. P. Bhattacharjee et al. / Scripta Materialia 56 (2007) 13–16
sintering temperature as a result of grain growth. The results indicate that change in low-energy boundary fraction during grain growth can be manifested only by an increase in low R (R3 boundaries in our case) boundary fractions. This observation can be correlated with the stacking fault energy value of pure Ni. An increase in the proportion of R3 boundaries is expected in Ni due to its lower stacking fault energy value and increased susceptibility to forming annealing twins compared to material with a higher stacking fault energy such as Al, in which the R3 proportion is found to be much lower [10]. Although the formation of annealing twins results in increased grain boundary area compared to the case in which no twins are formed, the total grain boundary energy can actually be reduced by the formation of annealing twins since the coherent twin boundary energy is much lower than random high-angle boundaries (typically 43 and 866 mJ m2 for Ni [11]). The criterion for formation and stability of a twin configuration from the point of view of overall energy reduction has been amply demonstrated by Burke [12] and Fullman and Fisher [13]. The observed increase in R3 fraction and simultaneous decrease in random boundary fraction with increasing sintering temperature indicate, therefore, that formation of annealing twins is a very favorable mechanism for reducing the overall grain boundary energy during grain growth. In conclusion: (1) The fraction of LAGBs (R1) does not change significantly with increasing sintering temperature.
(2) The fraction of random boundaries decreases and that of the R3 boundary increases with increasing sintering temperatures. (3) The increase in the R3 fraction is consistent with the observation of large numbers of twins at higher sintering temperatures. (4) The presence of a high fraction of R3 boundaries indicates that a grain boundary texture develops in pure Ni after sintering, although the materials do not possess any prominent crystallographic texture. [1] F.J. Humphreys, M. Hatherly, Recrystallization and Related Annealing Phenomena, Pergamon Press, Oxford, 1995. [2] H. Gleiter, B. Chalmers, Prog. Mater. Sci. 77 (1972) 16. [3] G. Palumbo, P.J. King, K.T. Aust, U. Erb, P.C. Lichtenberger, Scripta Metall. Mater. 1775 (1991) 25. [4] D. Dimos, P. Chaudhari, J. Mannhart, F.K. LeGoues, Phys. Rev. Lett. 219 (1988) 61. [5] J. Furley, V. Randle, Mater. Sci. Technol. 7 (1999) 12. [6] V. Randle, A. Brown, Philos. Magn. A 59 (1989) 1075. [7] Y. Pan, B.L. Adams, Scripta Metall. 30 (1994) 1055. [8] G.S. Grest et al., Acta Metall. 33 (1985) 509. [9] F.J. Humphreys, Scripta Metall. 27 (1992) 1557. [10] V. Randle, The Role of the Coincidence Site Lattice in Grain Boundary Engineering, The Institute of Materials, London, 1996. [11] L.E. Murr, Interfacial Phenomena in Metals and Alloys, Addison-Wesley, Reading, MA, 1975. [12] J.E. Burke, Trans. Metall. Soc. AIME 188 (1950) 1324. [13] R.E. Fullman, J.C. Fisher, J. Appl. Phys. 22 (1951) 1350.
Effect of sintering temperature on grain boundary character distribution in pure nickel P.P. Bhattacharjee, S.K. Sinha and A. Upadhyaya* Department of Materials and Metallurgical Engineering, Indian Institute of Technology, Kanpur 208 016, UP, India Received 25 May 2006; revised 29 August 2006; accepted 1 September 2006 Available online 10 October 2006
The evolution of grain boundaries in pure Ni has been studied after sintering at 900, 1100 and 1300 C. It is found that the fraction of low-angle grain boundaries remains more or less constant over this sintering temperature range. However, the fraction of random boundaries decreases and that of R3 boundaries increases with increasing sintering temperature. The observations on grain boundary character distribution are found to correlate well with the accompanying microstructural changes. 2006 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Nickel; Sintering; Grain growth; EBSD; Grain boundary
The phenomenon of grain growth following recrystallization has been well documented in the literature [1]. Grain growth has often been found to be associated with the development of a characteristic recrystallization texture and accompanying changes in grain boundary character distribution in a number of systems. The effect of grain growth and other annealing processes on grain boundary character distribution has been an area of significant interest due to the fact that grain boundaries can substantially influence properties such as creep resistance [2] and susceptibility to intergranular corrosion [3], and, more recently, critical current density in superconducting thin films has been found to be dependent on grain boundary character [4]. Enhancing the properties of materials by tailoring grain boundaries for a number of critical applications through ‘‘grain boundary engineering’’ has become an established and active area of materials research. However, grain growth and other annealing phenomena have generally been studied in great detail in materials that have been cold worked prior to the annealing treatment. Interestingly, in sintering, which is also a high-temperature treatment and is employed for densification of powder metallurgically prepared porous compacts, grain growth has been found to play a significant role. However, the evolution of grain boundaries * Corresponding author. Tel.: +91 512 2597672; fax: +91 512 2597505; e-mail: [email protected]
has been studied to a much lesser extent in these materials. This is surprising since powder metallurgically prepared materials are often of near net shape and hence there is a greater need to study the evolution of grain boundaries during sintering in such materials. With this in mind, the current work investigates the evolution of grain boundaries and the accompanying microstructural changes in pure Ni during sintering at different temperatures. Figure 1 shows a scanning electron microscope (SEM) micrograph while Table 1 summarizes the important characteristics of the Ni powder used in this study. It should be noted that the average particle diameter in Table 1 is evidently more than what appears in the micrograph of Figure 1. This may be because a wet measurement system (as in Table 1) usually gives larger particle diameters, due to particle agglomeration, compared to measurements on the population base (as in the SEM micrograph of Fig. 1). Pure Ni compacts were made using a rectangular die of dimensions 8 · 25 mm in a 50 ton hydraulic press (Model CTM 50, FIE, India) applying uniaxial pressure of 500 MPa. The green samples were sintered for 1 h at 900 C (±5C), 1100 C (±5 C) and 1300 C (±5 C), respectively. Sintering was carried out in a SiC-heated horizontal tubular furnace (rating 1.5 kV A) using commercially pure hydrogen as the sintering atmosphere (dew point 35 C). The heating rate (5 C min1) was kept constant for all the sintering operations. Green and sintered densities were determined from the dimensional
1359-6462/$ - see front matter 2006 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.scriptamat.2006.09.003
14
P. P. Bhattacharjee et al. / Scripta Materialia 56 (2007) 13–16
Figure 1. SEM micrograph of the Ni powder.
Table 1. Characteristics of the experimental Ni powder Powder Vendor Purity Shape Particle size (lm)
D10 D50 D90
Ni Sigma–Aldrich 99.99% Spiky 3.9 9.7 26
measurements of the green and sintered compacts, respectively, and the densification parameter was calculated from these values using the relation: SD GD ; TD GD where SD is sintered density, GD is green density and TD is theoretical density. Specimens for optical microscopy studies were prepared using usual metallography techniques. Samples for orientation imaging microscopy (OIM) studies were prepared using electropolishing at room temperature with platinum as cathode and the samples as anode. The OIM studies were conducted by a fully computer-controlled electron backscattered diffraction (EBSD) system attached to a SEM (QUANTA 200, FEI) using TSL OIM ANALYSIS Version 4.0 software. The step sizes were chosen such that at least 3–4 Kikuchi patterns were collected from each grain. For statistical accuracy, approximately 2000 grains were covered for each sample. Table 2 gives the sintered densities and densification parameters of pure Ni sintered at three different temperatures. The sintered density is found to increase in a consistent manner with increasing sintering temperatures. Figure 2(a)–(c) shows the optical micrographs of pure Ni sintered at 900, 1100 and 1300 C, respectively. Annealing twins are clearly visible in the microstructures Densification parameter ¼
Table 2. Effect of sintering temperature on the densification behavior of pure Ni Sintering temperature, C
Sintered density, % theoretical
Densification parameter
900 1100 1300
82.0 90.0 94.0
0.47 0.68 0.73
Figure 2. Optical micrographs of Ni compacts sintered at (a) 900 C, (b) 1100 C and (c) 1300 C.
of the materials sintered at 1100 and 1300 C as compared to the material sintered at 900 C. The micrographs of materials sintered at 1100 and 1300 C clearly indicate signs of grain growth, a finding that is also substantiated by the grain size distributions (Fig. 3) for these three sintering temperatures. The average grain size was around 6 lm after sintering at 900 C and around 11 and 35 lm after sintering at 1100 and 1300 C, respectively. The grain boundary character distributions (GBCDs) of pure Ni at these three sintering conditions reveal that the fraction of low-angle grain boundaries (LAGBs) is quite low after sintering at 900 C and does not change to any significant extent after sintering at 1100 and 1300 C. However, the coincidence site lattice (CSL) and random grain boundary fractions differ markedly in these three conditions. The fraction of random grain boundaries is quite high after sintering at 900 C and decreases slightly after sintering at 1100 C and quite significantly after sintering at 1300 C.
P. P. Bhattacharjee et al. / Scripta Materialia 56 (2007) 13–16
Figure 3. Grain size distribution in pure Ni after sintering at (a) 900 C, (b) 1100 C and (c) 1300 C.
15
boundaries in face-centered cubic (fcc) metals and alloys belong to the R3 (60/h1 1 1i relationship) category. The sintered microstructures and misorientation distribution plots also indicate that the R3 boundaries present can be identified exclusively as twin boundaries. The increase in R3 boundary fraction shows that the twin density increases with increasing sintering temperatures. In order to quantify the above observations, unique grain color maps were constructed from the EBSD datasets for these three sintered materials and volume fractions of the parent grains (P) or grains containing twins, the twinned regions within the parent grains (T) and the untwinned (U) grains (schematically shown in Fig. 5(a)) were determined (Fig. 5(b)) for the three sintering temperatures. Quantitative examination also shows clearly that the twin density increases with increasing sintering temperature. It has previously been suggested from experimental evidence [5–7] on different materials, as well as by computer simulations [8,9], that in a microstructure with a distribution of grain boundary energies the boundary tension would lead to a decrease in the total number of high-energy boundaries on annealing. In pure Ni, Furley and Randle [5] have reported an increase in R3 boundaries and a decrease in R5 boundaries during grain growth. However, in most cases the materials have been prepared through thermomechanical processing and as a result it may be difficult to pinpoint whether primary recrystallization or grain growth is the dominating effect in changing the GBCD. However, it can be seen that there is a general tendency for the proportion of low-energy grain boundaries to increase and high-energy grain boundaries to decrease during grain growth. Mobility of boundaries during grain growth has been thought to play a significant role, and a large increase is predicted for low-energy/low-mobility boundaries, such as LAGBs (R1) and other low R boundaries [9]. In the present study, pure Ni samples consolidated by the powder metallurgy technique have been investigated. Therefore, the complex effects of thermomechanical processing may be neglected and the change in GBCD can be solely attributed to changes taking place at the grain growth stage of sintering. The GBCDs clearly show that there is no significant change in LAGB proportion during grain growth. However, the fraction of random boundaries decreases and the proportion of R3 boundary increases with increasing
Figure 4. The CSL boundary distributions in pure Ni after sintering at (a) 900 C, (b) 1100 C and (c) 1300 C.
The CSL boundary fraction shows almost a reverse trend (Fig. 4(a)–(c)). The fraction of CSL boundary is rather low for the 900 C sintered sample and increases slightly after sintering at 1100 C and quite significantly after sintering at 1300 C. The CSL boundary distributions in these three conditions clearly indicate that most of the CSL boundaries belong to the R3 category. Boundaries beyond R3 are present in very minor proportions in all the sintered materials. Annealing twin
Figure 5. (a) Schematic representation of P, T and U (see text) and (b) histogram showing their respective volume percentages in the three sintered conditions.
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P. P. Bhattacharjee et al. / Scripta Materialia 56 (2007) 13–16
sintering temperature as a result of grain growth. The results indicate that change in low-energy boundary fraction during grain growth can be manifested only by an increase in low R (R3 boundaries in our case) boundary fractions. This observation can be correlated with the stacking fault energy value of pure Ni. An increase in the proportion of R3 boundaries is expected in Ni due to its lower stacking fault energy value and increased susceptibility to forming annealing twins compared to material with a higher stacking fault energy such as Al, in which the R3 proportion is found to be much lower [10]. Although the formation of annealing twins results in increased grain boundary area compared to the case in which no twins are formed, the total grain boundary energy can actually be reduced by the formation of annealing twins since the coherent twin boundary energy is much lower than random high-angle boundaries (typically 43 and 866 mJ m2 for Ni [11]). The criterion for formation and stability of a twin configuration from the point of view of overall energy reduction has been amply demonstrated by Burke [12] and Fullman and Fisher [13]. The observed increase in R3 fraction and simultaneous decrease in random boundary fraction with increasing sintering temperature indicate, therefore, that formation of annealing twins is a very favorable mechanism for reducing the overall grain boundary energy during grain growth. In conclusion: (1) The fraction of LAGBs (R1) does not change significantly with increasing sintering temperature.
(2) The fraction of random boundaries decreases and that of the R3 boundary increases with increasing sintering temperatures. (3) The increase in the R3 fraction is consistent with the observation of large numbers of twins at higher sintering temperatures. (4) The presence of a high fraction of R3 boundaries indicates that a grain boundary texture develops in pure Ni after sintering, although the materials do not possess any prominent crystallographic texture. [1] F.J. Humphreys, M. Hatherly, Recrystallization and Related Annealing Phenomena, Pergamon Press, Oxford, 1995. [2] H. Gleiter, B. Chalmers, Prog. Mater. Sci. 77 (1972) 16. [3] G. Palumbo, P.J. King, K.T. Aust, U. Erb, P.C. Lichtenberger, Scripta Metall. Mater. 1775 (1991) 25. [4] D. Dimos, P. Chaudhari, J. Mannhart, F.K. LeGoues, Phys. Rev. Lett. 219 (1988) 61. [5] J. Furley, V. Randle, Mater. Sci. Technol. 7 (1999) 12. [6] V. Randle, A. Brown, Philos. Magn. A 59 (1989) 1075. [7] Y. Pan, B.L. Adams, Scripta Metall. 30 (1994) 1055. [8] G.S. Grest et al., Acta Metall. 33 (1985) 509. [9] F.J. Humphreys, Scripta Metall. 27 (1992) 1557. [10] V. Randle, The Role of the Coincidence Site Lattice in Grain Boundary Engineering, The Institute of Materials, London, 1996. [11] L.E. Murr, Interfacial Phenomena in Metals and Alloys, Addison-Wesley, Reading, MA, 1975. [12] J.E. Burke, Trans. Metall. Soc. AIME 188 (1950) 1324. [13] R.E. Fullman, J.C. Fisher, J. Appl. Phys. 22 (1951) 1350.